Carbides precipitated from the melt in a Zr-2.5Nb alloy

Journal of Nuclear Materials 195 (1992) 265-276 North-Holland

Carbides precipitated

from the melt in a Zr-2SNb

alloy

R. Piotrkowski a, G. Vigna b, S.E. BermGdez b and E.A. Garcia a a Dto. Combustibles Nucleares, b Dto. Materiales, Comisih National de EnergLa Atdmica, Au. de1 Libertador 8250, 1429 Buenos Aires, Argentina

Received 21 April 1992; accepted 16 June 1992

An experimental method is presented which leads to the formation of carbides similar in size (3 to 8 km) and composition to those observed in pressure tubes of CANDU-type reactors. The method is based on melting of the Zr-2.5Nb alloy in a graphite crucible, where isothermal C diffusion in the Zr-Nb melt took place. As a result of the diffusion couple liquid Zr-2.5Nb/solid graphite, a carbide layer, up to 100 urn thick, grew attached to the crucible wall, together with carbide particles whose size was in the micrometer range. The smallest particles were arranged in rows determined by the prior P-phase grains. The main carbide phase detected was the cubic MC,_,; hexagonal M,C was also detected (M stands for metal). In some experiments that involved quenching from the p-phase, the w-phase was detected. Its occurrence was ascribed to the interstitial atoms (C, 0, N) present in the samples.

1. Introduction In a previous work [l], we studied the precipitation of carbides in a Zr-2.5Nb alloy. In that work, carbon diffusion experiments were performed. Precipitation of carbides occurred at 555”C, mainly in the hexagonal (hP2) a-(Zr, Nb) phase, and at 716°C in the c1(Zr, Nb) + P-(Zr, Nb) two-phase field, after heat treatment at 1100°C in the cubic (~12) l3-(Zr, Nb) phase [2]. The precipitated particles were identified as (Zr, Nb)C,_, by X-ray diffraction, scanning electron microscopy and electron microprobe analysis. (Zr, Nb)C,_, is a cubic (cFS) phase and it is hypostoichiometric, with x varying between 0.41 and 0.02 [3]. The semiquantitative result was (in atomic percent): C = 35, Nb = 2, Zr = 63, for the sample treated at 555°C comparing well with the Zr-C phase diagram [4], when Zr + Nb is considered and when the carbide phase is formed from the Zr-rich region (x = 0.41). The distribution of carbides was explained by C diffusion in the ol-(Zr, Nb) and p-(Zr, Nb) phases and in the o-l3 phase boundaries. The precipitation process was treated in terms of C diffusion in the o-(Zr, Nb) phase and this was in accordance with the size of the obtained particles with diameters ranging between 0.5 and 1.5 Frn. Hexagonal (hP9) M,C type carbides (one of the possibilities for Nb and (Zr, Nb) carbides [5,6]) were not detected in that work. Some indication Elsevier Science Publishers B.V.

from the X-ray diffraction patterns suggested that (Zr, Nb)C,_, particles could have grown with internal stresses, although the slight shifts in diffraction peak positions could also be due to quenching effects. Carbides can play a deleterious role in pressure tubes of CANDU reactors. One of the possible reasons for the failure of pressure tubes, before the predicted lifetime has expired, is the hydrogen induced delayed cracking. The presence of surface carbides could affect the protective oxide layer integrity, and hence drastically augment the hydrogen pick up. This would contribute to fast fracture of tubes, and then to the break before leak condition. Moreover, if the particles are sufficiently large (say diameter > 1 km), they could create around them a deformation field that could modify the H solubility and in this way promote the hydrogen pickup enhancement. The purpose of the present work is to search for a method to obtain carbides similar in size and composition to those observed in a pressure tube material (P.T.), which will help to identify the stage in the fabrication route at which these carbides originate.

2. Experimental The Zr-2.5Nb alloy was prepared by Wah Chang as cold-rolled and stress-relieved 1 mm thick strips. From

R. Piotrkowski et al. / Carbides from the melt in a Zr-2.5Nb

266

alloy

these, several samples were prepared in different ways, with the purpose of obtaining carbides with the same size as that observed in samples coming from a P.T., before its service in a nuclear reactor. These P.T. carbides were polyhedral particles with a typical size of some micrometers and poor in Nb. The obtainment of the described carbides was attempted by three methods: (1) carbon diffusion and carbide precipitation after evaporating graphite on one polished face of each sample; (2) carbon diffusion in a carbon rich gaseous environment; (3) carbon diffusion in the Zr-2.5Nb melt.

After diffusion annealing of sample 2 and after precipitation annealing of sample 1, they were sectioned through a transversal plane. This transversal face was abraded with emery paper down to 600 grit, then polished with diamond paste of 7 and 3 pm and finally chemically etched with a 45% glycerin-45% HNO,-10% HF solution, and observed by scanning electron microscopy @EM). X-ray diffraction patterns were obtained from the top and the bottom face of each sample (top was where C had been deposited) with a horizontal diffractometer using Ko Co radiation.

2.1. Method 1

2.2. Method 2

Carbides were previously precipitated by this method in ref. [l]. The procedure starts with vacuum evaporation of graphite of spectral purity on one polished face of each sample. The thickness of the carbon layer was evaluated using the weight difference and the normal graphite density (.p = 2.26 Mg/m3). The annealing treatments were performed in ChevenardJoumier furnaces with AT = f 1°C. The samples were wrapped in high-purity Ta foils and vacuum-sealed in quartz tubes with a slight overpressure of high-purity Ar. The thickness of the carbon layer and the duration of the diffusion annealing were modified in this work with respect to ref. [l] in order to obtain a uniform distribution of carbides through the samples thickness (1 mm). The duration of the precipitation annealing was modified in order to obtain particles with a similar size to that observed in the sample from the pressure tube. Table 1 shows the preparation route for sample 1 (submitted to diffusion and precipitation annealings) and sample 2 (submitted only to diffusion annealing). After the diffusion annealings, if the total C deposit entered the samples, and no evaporation occurred, the C concentration would be 0.01 wt% for both samples, according to eq. (1) in ref. [l]. After this procedure, and according to eq. (2) in ref. [I], carbides with a mean diameter of = 3 km were expected in sample 1.

Samples were submitted to a heat treatment at 930°C during 3 h in a dynamic atmosphere that consisted of a mixture of methane, propane and air, widely used to carburize steels, and finally cooled inside the furnace, in the same atmosphere. After heat treatment a transversal face was prepared for observation by SEM, in the same conditions described under method 1. X-ray diffraction patterns were obtained (e.g. sample 3). Total C content measurements were performed in several samples with the method consisting in the conversion to CO, in an induction furnace and detection by absorption of infrared radiation.

Table 1 Preparation route for samples by method 1. Both samples were quenched in water Sample

Weight (8)

C deposit (g)

Diffusion

Precipitation

CC)

(h)

(“Cl

(h)

1 2

1.48050 1.88599

1.5~10-~ 2.1~10-~

1080 1125

17.7 17

549 -

1439.5 _

2.3. Method 3 In a first experiment 0.64 g of the Zr-2.5Nb alloy (sample 4) was set in a cylindrical graphite crucible with 11 mm inner diameter, 14 mm outer diameter and 14 mm height, and melted in a Brew furnace under high vacuum conditions (diffusing pump conditions). The crucible positioned in the center of the furnace chamber was held by a basket made of Ta. The temperature was measured with a W-SRe/W-26Re thermocouple touching the outer surface of the bottom of the Ta basket. Due to the fact that contamination with Ta took place, a second experiment was performed. 5.4 g of the Zr-2.5Nb alloy (sample 5) was melted in a covered cylindrical crucible with 14 mm inner diameter, 20 mm outer diameter and 35 mm height. The cap was a graphite disk with a 4 mm diameter central hole. Table 2 indicates the annealing conditions for samples 4 and 5. Then, for each sample, the set crucible plus melted alloy was cut through a plane perpendicular to the crucible axis. After abrading with emery papers down

R. Piotrkowski et al. / Carbides from the melt in a Zr-2.5Nb Table 2 Preparation

267

route for samples by method 3

Sample Mean heating rate

4 5

allov

Melting conditions

pCmin_,)

CT)

46 55

1870+30 60 1850f 10 24

Mean cooling

Pressure (bar)

60 66

<5x10-s <3x10-s

to 600 grit and polishing with diamond paste down to 1 urn, the sample was studied with a CAMECA SX 50 microprobe equipped with a wavelength dispersive system. An X-ray diffraction pattern was obtained from sample 5.

3. Results and discussion 3.1. Method

1

Fig. 1 shows two SEM micrographs of sample 1, obtained after different chemical etchings. Very few and very small carbide particles could be observed. The dark phase is a-Widmanstltten, randomly oriented. The bright particles are mainly cw’-phase, prior l3, enriched in Nb. Some of them are carbides. This result was corroborated by X-ray diffraction. The trace of many extracted hydrides can be seen in fig. la where the polished face was chemically etched with a 45% glycerin-45% HNO,-10% HF solution. Some hydrides are shown in fig. lb), where the magnification is higher. In this case the chemical etching was performed using a 45% lactic acid-45% HNO,-10% HF mixture. In tables 3 and 4 the X-ray diffraction patterns obtained in the studies of the front (F) and back (B) faces of samples 1 and 2 are compared with data from the literature [7]. It can be seen that the phases detected in sample 1 are cw-Zr and ZrC. However, in comparison with the diffraction pattern of the sample treated at 555°C in ref. [l], where both the diffusion annealing and the precipitation annealing were markedly shorter, the intensity corresponding to ZrC in the present study was much lower. The lattice parameters of the phases obtained by Cohen’s method of least squares were, for o-phase, from F and B respectively, are a = (0.3236 f 0.0002) nm, c = (0.5165 + 0.0003) nm, c/a

= (1.596 + 0.002);

Fig. 1. SEM micrographs of sample 1. The dark phase is u-Widmanstltten, randomly oriented. The bright particles are mainly a’-phase, prior p, enriched in Nb. Some of them are carbides. (a) After 45% glycerin-45% HNO,-10% HF etching, hydrides were extracted. (b) After 45% lactic acid-45% HNO,-10% HF etching, hydrides were revealed.

a = (0.3233 + 0.0002) nm, c = (0.5161 k 0.0003) nm, c/a

= (1.596

f 0.002) ;

and for the carbide phase a = (0.4713 + 0.0002) nm. The o-phase parameters compare well with the literature data [2], that are u = 0.3232 nm; c = 0.5147 nm. The carbide parameter compares quite well taking into account that we had only three peaks, and with low intensity. In ref. [3], for ZrC,, s5 the parameter a = 4.691 nm. There arc two reasons why the predicted size of carbide particles did not occur. First, the available C mass was less than foreseen. A loss of material during the diffusion annealing took place; to check this assertion, the samples were weighed after treatment (sample 1 after diffusion plus precipitation anneals and sample 2 after diffusion anneal). Table 5 summarizes these weight changes. For both samples the weight loss

268

R. Piotrkowski et al. / Carbides from the melt in a Zr-2SNb

a = (0.3234 + 0.0002) nm, c = (0.5166 f 0.0003) nm, = (1 S97 f 0.002). The rest of the obtained reflections lie between the p and the w lines, indicating the presence of the diffuse w-phase, typical in Zr alloys, as is described below. So, it makes not sense to calculate the lattice parameters from our data in this case. The o-phase (hexagonal (hP3)) occurs in pure Zr at high pressures. Previous work in the literature, summarized in refs. 19,101 deal with the occurrence of the w-phase in Zr alloys with other d-rich transition elements. In alloys at atmospheric pressure, the o-phase forms as clusters imbedded in the bee matrix. These o-domains are lo-20 nm in diameter. By rapid quenching to room temperature from the P-phase in Zr-Nb alloys two metastable phases occur without change of composition, besides the retained P-phase: (Y’ and w. The al-phase is a close-packed hexagonal phase that forms athermally during rapid quenching from the P-phase to room temperature in Zr-Nb alloys with less than 7.5 wt% Nb [2]. It has the crystal structure of the a-phase and the composition of

was of the order of the nominal oxygen content. During the long-time diffusion annealing, besides the C homogenization, the sample deoxidation took place in the presence of the C atoms thus producing CO or CO,. On the other hand, as the dissolved 0 enlarges the H solubility [8], the deoxidation produced the precipitation of hydrides. Secondly, since in the present work C was uniformly distributed, the carbides precipitated near the top and bottom faces would in principle be less than in the previous work and this would enhance the effect of diminishing the carbide signal in the X-ray diffraction patterns. The Co Ka X-ray diffraction gives information from a slab = 7 urn from the surface. According to table 4, no carbide phases were found in sample 2, a result which agrees with the accepted Zr-C phase diagram [4]. In this sample the detected phases were a-Zr, p-Zr and o-Zr. The o-phase lattice parameters, from F and B were respectively:

c/a

u = (0.3231+ 0.0002) nm, c = (0.5161 + 0.0003) nm, c/a

=

alloy

(1 S97 * 0.002);

Table 3 X-ray diffraction pattern of sample 1 Front face F

Back face B

ci-Zr

ZrC

d,

IF

‘f,

d

I

(pm)

(asi.)

(pm)

(pm)

(a.u.)

280.3 271.1 258.3 247.3 234.8 190.1 166.5 162.1 147.1 140.2 137.1 135.2 128.9 123.2 117.1 108.8 106.0 103.9 100.9

30 2 15 100
281.0

13

279.8

33

100

258.3 247.3

4 100

257.3 245.9

32 100

002 101

190.1


189.4

17

102

162.1 146.9 140.4 137.2

9 100 1 2
4

129.3 122.9 117.0 108.8 105.9 103.7 100.7 97.8 96.9

4

93.3

161.6 146.3 139.9 136.8 135.0 128.7 123.0 116.9 108.4 105.9 103.6 100.6 97.8 96.6 94.7 93.3

17 18 3 18 12 4 4 3 4 2 6 3 2 4 2 2

110 103 200 112 201 004 202 104 203 210 211 114 212 105 204 300

97.1 94.8 93.3

2 4

hkl

hkl

d

(pm)

fax.)

270.9

100

111

234.6

82

200

165.9

62

220

R. Piotrkowski et al. / Carbides .from the melt in a Zr-2.5ti

269

alloy

Table 4 X-ray diffraction pattern of sample 2 Front face F d, (pm) 281.3 258.9

247.0 198.8 190.6

11 11

100 1 4

161.7 146.7 144.8

22 89 22

137.2 128.9 123.0 117.1

9 5 6 4

105.7 103.6 100.9

16 81 67

Back face B

a-Zr

da (pm)

d (pm) 33 32

12 16 5

279.8 257.3

247.1 198.9 190.0 183.2 169.4 161.7 146.8 144.9

44 1 13 1 1 59 100 14

245.9

100

101

189.4

17

102

137.3 129.1

1 5

117.2 108.6 105.9 103.6 101.0

38 12 17 66 33

136.8 128.7 123.0 116.9 108.4 105.9 103.5 100.6

17 18

18 4 4 3 4 2 6 3

the prior P-phase. The o-phase in alloys is a variant of the hexagonal (hP3) o-phase occurring in Zr at high pressures. The crystal structure depends on composition. It forms in Zr-Nb alloys with more than 7 wt% Nb, coexisting between 7 and 7.5% with (Y’. After the 17 h anneal at 1125”C, the only phase should be p-Zr. In our quenching experiment, the Nb content (- 2.5 wt%,) was not in the range that corresponds to the p-w transformation. The C content (less than 0.01 wt%) seems to have been sufficient to promote it. The p-o transformation was also detected in our previous work [l], where it was also ascribed to the presence of the C atoms. According to different electronic structure models of the o-phase in Zr [lo], covalent bonds based on sd* hybridised orbitals occur,

Table 5 Evolution of weight in samples 1 and 2 Sample

1 2

Initial weight

Weight with C deposit

Weight after heat treatment

(g)

(g)

(g)

1.48050 1.88599

1.48065 1.88620

1.47929 1.88528

hkl

d (pm)

[ax.)

250.7

100

hkl

d (pm)

100 002

281.4 258.1 251.3

161.6 146.3

w-Zr

p-Zr hkl

110

252.7

100

101 110

196.1

10

111

177.3

17

200

179.2 156.6

40 20

201 002

144.7

33

211

145.7

80

121 300

110 103

112 004 202 104 203 210 211 114

similar to the sp’ hybridisation in the graphite structure. The interstitial C-atoms could contribute to establish similar hybridised orbitals in our Zr alloy. The same could be valid for other electron donor interstitial atoms in transition metals. In ref. [ll] the authors demonstrated by elastic-diffuse neutron scattering on P-Nb samples with very low content of O-interstitial atoms, that the lattice distortions around a single interstitial 0 impurity give place to o-phase embryos. 3.2. Method 2 Method 2 failed in producing carbides, but an important oxide scale was developed on the sample in a typical oxidation process. Fig. 2 shows SEM micrographs of the sample. Four zones were obtained which are denoted as A, B, C and D. A is the oxide layer, 15 km thick. B is a two-phase zone, 35 urn thick, where the main phase resembles a-(Zr, Nb) obtained in previous work [1,12,13], which in the present case would have been favored by oxygen contamination. This is suggested by the Zr-0 [14] and the Zr-Nb-0 phase diagrams [ 151.A second dispersed phase rich in Nb can be seen. This is probably a prior metastable P-phase

270

R. Piotrkowski et al. / Carbides from the melt in a Zr-2.5Nb

which was also found in our previous work [l], and by other authors in a Zr-2SNb alloy [16]. This phase transformed to (Y’ by quenching. The C zone, 120 km thick, is the prior (a + l3) phase structure, which may also be taken as a sign of the oxygen contamination, since in the binary Zr-Nb phase diagram [2] an alloy with 2.5% Nb at 930°C consists only of a single P-phase. Region D, the central zone, is the prior P-phase. During the cooling process, that took place in the furnace, the a-phase has preferentially nucleated and grown around the prior /3 grain borders, as was described in [ 12,131. This process was also enhanced by the presence of oxygen. Tables 6 and 7, respectively, show the X-ray diffraction pattern for sample 3 before and after extracting the oxide scale and compare with data from the literature [7]. In table 6 the more intense peaks come from the monoclinic (mP12) ZrO,; there are diffraction peaks from tetragonal (tP6) ZrO,, stable at temperatures higher than 1200°C in the Zr-0 system, stabi-

Table 6 X-ray diffraction Experiment d

pattern

of sample Monoclinic

3 before

extracting

hkl

I (a.u.)

d

(pm)

(pm)

(a.u.1

371.2 318.2 297.2 286.7 264.7 252.9

10 100 18 7 61 29

369.0 316.0

18 100

011 11

283.4 261.7 253.8

65 20 14

11 002 200

222.2 200.2 185.5 181.4 166.4

12 34 11 15 31

221.3 201.5 184.5 181.8 165.6

14 8 18 12 14

21i 112 022 220 013 221

161.7

6

160.8

8

154.6 150.5 147.9

23 23 6

154.1 150.8 147.6

10 6 6

3ii 212 131 113 311

145.6

4

144.7

4

121

137.8

2

133.1 122.7 118.4

25 3 5

132.1

6

lized in this case by the other interstitial atoms (C, N), and diffraction peaks that can be ascribed to the hexagonal o-phase. The lattice parameters obtained from table 6 are, for the monoclinic oxide phase: a = (0.516 f 0.001) nm, b = (0.521 f 0.002) nm. c = (0.532 f 0.002) nm, j3 = (99.5 + 0.3)“, a/b c/b

= (0.991 + 0.002),

= (1.021 f 0.002);

for the primitive

tetragonal

oxide phase:

a = (0.362 f 0.001) nm, c = (0.521 + 0.001) nm; and for the hexagonal

w-phase:

a = (0.504 f 0.001) nm, c = (0.313 rf: 0.001) nm, c/a

= (0.621 + 0.001).

For the three phases the parameters compare well with those from literature. From ref. [14] the parameters are, for the monoclinic oxide a = 0.5169 nm, b = 0.5232

the oxide scale Tetragonal

ZrO,

alloy

w-Zr

ZrO, hkl

d

(pm)

:a.u.l

296.0

100

111

260.0 254.0

18 25

002 200

181.0

35

220

154.7

45

131

d

hkl

(pm)

:a.“.)

252.1

100

101 110

145.7

80

121 300

133.0

30

112

1

137.7

2

132 123

123.3 119.0

2 12

114 133

104

R. Piotrkowski et al. / Carbides from the melt in a Zr-2.5Nb Table 7 X-ray diffraction pattern of sample 3, after extracting the oxide scale Experiment

p-Zr

a-Zr

d

I

d

I

(pm)

(au.1

(pm)

(a.u.1

hkl

281.0 258.5 247.8 190.5 177.5 162.0 146.9 144.2 140.3 137.1 135.6 129.2 117.4 108.86 106.00 103.90 loo.97 98.10 96.91

7 27 38 21 6 54 68 7 < 1 26 1 4 4 3 5 15 100 2 32

279.8 257.3 245.9 189.4

33 32 100 17

100 002 101 102

161.6 146.3

17 18

110 103

139.9 136.8 135.0 128.7 116.89 108.42 105.88 103.60 100.63 97.83 96.60

3 18 12 4 3 4 2 6 3 2 4

200 112 201 004 104 203 210 211 114 212 105

d

I

(pm)

(a.u.)

hkl

alloy

271

Table 7 shows the diffraction pattern after extraction of the oxide layer. The only phases detected were o-Zr and @Zr. No carbide phases were detected. The lattice parameters were, for the a-phase: a = (0.3240 _t 0.0001) nm, c = (0.5166 + 0.0001) nm, c/a = (1.594 * 0.001);

177.3

17

200

144.7

33

211

nm, c = 0.5341 nm, fi = 99,25” , and for the tetragonal oxide a = 0.35882 nm, c = 0.51882 nm. From ref. [2] the parameters of the o-phase are a = 0.5039 nm, c = 0.3136 nm. According to the assessed Zr-Nb-0 phase diagram at 1000°C (fig. 11 in ref. [15]), and considering the simplified scheme of a straight line as diffusional path, between the Zr-2.5Nb point and the 0 corner, our oxide layer can be expected to have various two-phase structures, such as (ZrO, + NbO,), (ZrO, + NbO), (ZrO, + P(Nb)) in which the P-phase is rich in Nb. The equilibrium @-Nb) would have a composition (in at%) ranging between = (96 Nb, 4 0) and (96 Nb, 4 Zr). Although in our case the equilibrium composition could have not been attained during the (930°C 3 h) heat treatment, both the Nb and 0 content could have been sufficiently high to promote the /3-w diffusionless transformation under cooling. The Nb oxides were not detected in our X-ray diffraction patterns. Although a systematic work on the influence of 0 (and other interstitials) in the promotion of the p-w transformation in Zr alloys is still lacking in the literature, there is some work that points at that direction; e.g. the p-o transformation induced by interstitial 0 and N in a Zr-Ti alloy containing 0.3 wt% 0 and 0.8 wt% N was detected in ref. [17].

and for the only two existing peaks the lattice parameter for the P-phase is: a = 0.3518 nm. From ref. [2], for the P-phase and 2.5% Nb, a = 0.3575 nm. The total C-content results were well within the nominal C-content range in commercial alloys. 3.3. Method 3 Melting of the alloy took place by method 3, C atoms entered the liquid alloy, and precipitation of carbides occurred. In sample 4, the melt climbed up the inner wall of the crucible and went down the outer wall, reaching the bottom of the Ta basket. This produced contamination of the sample because Ta was able to diffuse through the liquid. Figs. 3 and 4 show SEM micrographs of different magnification of the alloys obtained from the melt that resulted in a layer attached to the crucible wall in sample 4 and 5, respectively. Two zones could be detected. The right zone in fig. 3a, touching the crucible wall, shows the carbide phase. Going towards the left we see the metallic matrix with rows of carbides, some of them with polyhedral shape. The underlying two-phase structure in the metallic region is faintly revealed. Table 8 shows the X-ray diffraction pattern from a plane perpendicular to the crucible axis in the middle of sample 5 and it is compared with data from the literature [7]. Four phases were detected: a-Zr, ZrC, Nb,C and graphite. The Miller indices ascribed to the M,C phase were according to the hexagonal carbide reported in ref. [5], although our diffraction peaks could also be ascribed to an orthorhombic carbide reported in card 19-858 in ref. 171.The lattice parameters obtained were for a-Zr: a = (0.3235 + 0.0001) nm, c = (0.5153 k 0.0001) nm; for ZrC: a = (0.4684 f 0.0001) nm; and for Nb,C: a = (0.539 + 0.002) nm, c = (0.4981 f 0.0003) nm. Comparing with the data from literature, for ZrC,,,, [3] a = 0.46983; for Nb,C [5] a = 0.5407 nm; c = 0.4968 nm. Chemical microanalysis results are given in fig. 5 and table 9. Fig. 5 shows the concentration profiles of

272

R. Piotrkowski et al. / Carbides from the melt in a Zr-2.5Nb

alloy

the different elements forming a polyhedral MC,_, carbide in sample 4. The carbide particle (it was a general result) is poor in Nb and Ta. The dark gray phase in the metal region is a-phase, poor in Nb and Ta, whereas the light gray phase is P-phase, richer in these elements. Carbides are generally surrounded by a-phase. Table 9 shows the results of the microanalysis measurements performed in the carbide phases in samples 4 and 5. For comparison, results from a P.T. material, before service are also presented. The Zr and C contents in our samples agree within the experimental error with the corresponding contents in the P.T. sam-

Fig. 3. SEM micrographs from sample 4. (a) On the right side the carbide layer in touch with the crucible wall can be seen. Going to the left we see the metallic matrix with a number of carbide particles, some of them arranged in chains. (b) A large carbide phase is on the bottom, a row of particles are on the top and the metallic two-phase structure is faintly revealed.

Fig. 2. SEM micrographs of sample 3. Four zones are noted as A, B, C in (a) and D in (b). A is the oxide layer, 15 pm thick. B is a two-phase zone, 3.5 km thick. The dark phase is &Zr, Nb). The second dispersed phase, rich in Nb, transformed to (Y’ by quenching. C, 120 km thick, is the prior (a + B)-phase structure. D, the central zone, is the prior P-phase.

ple. The Nb (or Nb + Ta) content is very low in all the samples, but is still lower in the P.T. sample, where it is close to the microprobe detection limit. The results of microprobe measurements on the MC ,_X phase compare well with the C-rich boundary of the carbide in the Zr-C phase diagram [4], when M = (Zr + Nb + Ta) is taken instead of Zr. In this work some very small M,C particles, rich in Nb and Ta, and with diameters smaller than 1 km were detected in sample 4. Fig. 6 shows a backscattered electron image of a small particle that was a M,C carbide. No M,C particles could be detected with microprobe analysis in sample 5 although signals of this phase were

R. Piotrkowski et al. / Carbides from the melt in a Zr-2.5Nb

273

alloy

the carbide layer, while the smaller ones are arranged in rows that resemble the P-grain borders obtained in the first stages of the solidification process, occurring slowly inside the furnace. These smaller particles could probably have been dragged to the interdendritic regions and follow the P-grains evolution during the higher temperature stages of the cooling process. This central microstructure is very similar to that obtained in ref. [19] after annealing of a Zr-2 at% C alloy at 1865”C, where the corresponding micrograph reveals chains or rows of primary ZrC grains. If we consider only the binary Zr-C phase diagram, we can guess that the carbide layer could first have nucleated after the diffusion of Zr atoms into the graphite in the first stage of the high temperature treatment. This would explain that the values of the chemical composition obtained from the MC,_, particles can be reached from the C-rich zone of the Zr-C phase diagram. Then the carbide phase could have grown by C (and/or Zr) diffusion through the threephase system: graphite, carbide, (Zr, Nb) alloy, in a typical scheme described in other work [20,21], where the second Fick law holds in each phase and the Stefan law governs the phase boundary movement at the

detected by X-ray diffraction. They could be very small particles, and perhaps associated with the larger MC 1ox particles. Zr only forms the cubic MC,_, carbide with the (cF8), NaCl-type structure, where M stands for metal [4]. Ta and Nb are isomorphous elements and are completely miscible in the whole composition range and behave in a similar way when alloyed with Zr [2,4]. They form two types of carbides, the cubic MC,_, carbide and the hexagonal (hP9) M,C carbide 1431. The results obtained from samples 4 and 5 were quite similar, verifying that Ta and Nb behave in a very similar way in the studied system. The carbide phase grew mainly between the graphite and the molten (Zr, Nb) alloy, but a number of carbide particles also appears in a scheme that corresponds to isothermal diffusion in ternary systems, where the same chemical composition and a gradient in the volume concentration in the diffusion direction are expected for the dispersed phase, in our case the carbide phase [18]. This occurred in fact in our work, where the chemical composition of the carbide layer and the carbide particles agrees within the experimental error. The larger carbide particles look like pulled out from

Table 8 X-ray diffraction pattern of sample 5 Experiment

a-Zr

d

d

(pm)

:a.“.)

338.2 280.1 270.5 257.8 249.1 246.2 234.7 213.7 203.6 189.7 168.2 165.7 161.8 146.5 141.3 140.0 137.0 135.2 128.8 124.6 123.1

85 18 8 64 22 100 5 2 3 1.9 25 3 17 4 2 2 8 6 5

zrc hkl

(pm)

fa.u.1

279.8

33

100

257.3

32

002

245.9

100

101 234.6

189.4

17

d

hkl

(pm)

100

82

111

270.3

110

248.4

002

d

hkl

(pm)

la.u.1

336.0

100

002

213.0 203.0

10 50

100 101

167.8

80

004

123.2

30

110

200

102

161.6 146.3

17 18

110 103

139.9 136.8

3 18

200 112

128.7

4

004

123.0

4

202

L

4

hkl

(pm)

270.9

C

%C

d

165.9

62

220

141.5

50

311

135.5

19

222

141.2

113

124.2

004

274

R. Piotrkowski et al. / Carbides from the melt in a Zr-2SNb

phase boundaries. Moreover, the liquid (Zr, Nb)-solid (Zr, NbK, pi interaction occurring at high temperatures seems to lead to the carbide dissolution as was described in ref. [22] for the high temperature liquid Zircaloy-solid uranium oxide interaction. According to this scheme, the rows of carbides observed in the micrographs in the middle of the metallic regions would be originated from carbide particles pulled out from the carbide layer and dragged to the interdendritic regions during the solidification process. If we consider the ternary Zr-Nb-C system, from the only isothermal section available at 1500°C [6], it can bc observed that the presence of Nb shifts the carbide composition to higher C contents. Due to this fact, the carbide composition obtained by us can be reached from the metal rich region of the ternary phase diagram. More experimental work is necessary

alloy

to measure the velocity of the interphase movement, and to verify the assumptions for a diffusion model.

4. Conclusions Method 1 proved to be able to produce a high density of carbides at 555°C with diameters near 1 urn in a nearly 100 (*rn thick region [l], however, the method could not be improved in this work in the sense of obtaining larger carbides uniformly distributed in the 1 mm sample width. During the long-time annealing at high temperatures destined to diffuse the C atoms through the sample width, C reacted with 0 present in the samples, thus producing deoxidation. The deoxidation of the samples promoted the precipitation of hydrides, since there is a relation between 0 content and H solubility. In the sample where C was uniformly diffused we detected the diffuse o-phase after quenching the sample from the P-phase. This phase is not expected to occur by quenching a Zr-2.5 Nb alloy from the P-phase, so its presence was ascribed to the influence of the interstitial C atoms. Method 2, annealing at 930°C in a methane-propane-air atmosphere and widely used to carburize steels, failed to carburize our Zr-2.5Nb ahoy. Thus, a carburizing atmosphere for Fe resulted in an oxidizing atmosphere for Zr. A display of different metallographic structures could be appreciated through the sample width according to the 0 gradient. Two types of oxides were detected by X-ray diffraction: the monoclinic oxide, stable down to room temperature in the Zr-0 system, and the tetragonal oxide, that only appears in the binary Zr-0 system above lOOo”C, and could in this case be stabilized due to the presence of C and N atoms. Also the o-Zr phase (prior P-phase) was detected; its occurrence can be explained by the influence of various interstitial atoms. Entering the

Table 9 Microanalysis results of carbides in samples 4 and 5. The composition of the carbide phase in pressure tube material (P.T.) is for comparison. Fluctuation in C determination was 10%

Fig. 4. SEM micrographs from sample 5. In (a) the carbide layer is on the right side, and in (b) more details can be seen on the bottom side.

Sample

Phase

Zr (at%,)

Nb (at%,)

Ta (at%)

C (at%)

Number particles

4 4 5 P.T.

MC M,C MC MC

52.1 55.0 52.7 52.4

0.2 2.5 0.2 0.02

0.1 6.4

47.6 36.0 41.1 47.6

10 2 44 23

of

R. Piotrkowski et al. / Carbides from the melt in a Zr-2.5M

alloy

Gg. 5. Concentration profiles of the different elements fixming a polyhedral carbide in sample 4. The dark gray phan se is is P-phase, richer in these elements. (a) C prc Ifile, :-phase, poor in Nb and Ta, whereas the light gra y phase (b) Zr profile, (c) Nb profile, (d) Ta profile.

sample, different two-phase structures appeared, all of them strongly influenced by oxygen. Method 3 resulted successfully in obtaining carbides similar in size and composition to those observed in some P.T. Both in pressure tube material and in our alloy obtained from the melt, the cubic carbide MC, _-x particles were very poor in Nb. This would indicate that carbides present in the P.T. could have been originated from high-temperature stages of the tube elaboration (e.g. elaboration of the arc-melted ingot). Otherwise they could already be incorporated in the zirconium sponge. The hexagonal M,C carbide phase was also detected in this work. The microanalysis results for the MC,_, phase compare well with the accepted Zr-C phase diagram, when M (metal content) is taken instead of Zr and when the carbide phase is formed from the C-rich

Fig. 6. Backscattered

electron image of a small M,C in sample 4.

carbide

276

R. Piotrkowski

et al. / Carbides from the melt in a Zr-2.5Nb

region of the Zr-C phase diagram. Certainly, the results should be compared with the ternary Zr-Nb-C phase diagram, where the carbide phase could have been formed from the metal-rich region. The carbide phase was formed mainly as a solid layer between the graphite crucible and the liquid metal alloy, controlled by C and/or Zr diffusion, although some dissolution seems to have occurred at the liquid metal-solid carbide interface. A number of carbide particles in the micrometer range can be seen in the middle of the alloy. The smaller carbide particles are arranged in rows that resemble the P-grain borders originating from the first stages of the solidification process.

Acknowledgements

This work was performed under the Contract N.6249 with the International Atomic Energy Agency. The authors wish to thank M. Maldovan and D. Hermida for their help in the obtention of the X-ray diffraction patterns, M.R. Eppis for the total C-content determinations, and Drs. R. Baggio (Dto. Fisica), E. Vicente and A. Fernandez Guillermet (Centro Atomic0 Bariloche) for very fruitful discussions.

References

t11 R. Piotrkowski,

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